Welding Coating.pdf

WELDING AND COATING METALLURGY
1 INTRODUCTION 4
1.1 ROLE OF CARBON IN STEEL 4
1.2 WELDABILITY OF STEELS 5
2 CRYSTAL STRUCTURE 7
2.1 SOLUBILITY OF CARBON 8
3 IRON - IRON CARBIDE PHASE DIAGRAM 9
3.1 AUSTENITE (γ) 10
3.2 FERRITE (α) 11
3.3 PERITECTIC 11
3.4 PEARLITE 12
3.4.1 PEARLITE GROWTH 13
3.5 PRO-EUTECTOID FERRITE 14
3.6 PHASE TRANSFORMATIONS IN LOW ALLOY STEELS 15
3.7 GRAIN GROWTH 16
3.8 NON-EQUILIBRIUM COOLING 17
3.9 MARTENSITE - EFFECT OF RAPID COOLING 18
3.10 BAINITE 19
4 TRANSFORMATION DIAGRAMS 20
4.1 TIME TEMPERATURE TRANSFORMATION (TTT) DIAGRAMS 20
4.2 CONTINUOUS COOLING TRANSFORMATION (CCT) DIAGRAMS 23
4.2.1 CRITICAL COOLING RATES 24
4.2.2 DETERMINING CCT DIAGRAMS 24
WELDING AND COATING METALLURGY2 12 October 1999

FOR INTERNAL USE ONLY
4.3 EFFECT OF ALLOYING ELEMENTS 24
4.4 Ms and Mf TEMPERATURES 25
5 HARDENABILITY / WELDABILITY OF STEELS 26
5.1 CARBON EQUIVALENT (CE) & WELDABILITY 28
5.2 TEMPERING – EFFECTS OF REHEATING 29
5.3 SECONDARY HARDENING 30
6 HYDROGEN CRACKING RELATED TO WELDABILITY 32
6.1 LAMELLAR TEARING 34
7 REHEAT CRACKING IN THE HAZ 36
8 WELDING STEELS CONSIDERED DIFFICULT 38
8.1 PROCEDURAL CONSIDERATIONS 38
8.2 POST WELD HEAT TREATMENT (PWHT) 38
8.3 THE HEAT AFFECTED ZONE (HAZ) 38
8.3.1 LOSS OF TOUGHNESS IN THE HAZ 41
8.4 PREHEAT & CARBON EQUIVALENT 41
8.4.1 SEFERIAN GRAPH 42
9 SUMMARY 44
10 GENERAL ASPECTS CONCERNED WITH WEAR 45
PROTECTIVE COATINGS 45
11 SELECTING THE OPTIMUM WEAR RESISTANT SOLUTION 46
12 METHODS OF DEPOSITION 47
13 WELDING PROCEDURAL GUIDELINES 49
WELDING AND COATING METALLURGY2 12 October 1999 2 of 69

13.1 BASE METAL CONSIDERATIONS 49
13.1.1 WELDABILITY FACTORS 49
14 APPLICATION OF WEAR PROTECTIVE COATINGS 50
14.1 BASE METAL PREPARATION 50
14.2 PREHEAT 50
14.3 BUILD-UP 51
14.4 APPLICATION TECHNIQUE 51
14.5 COOLING PROCEDURE 51
14.6 FINISHING 51
15 WEAR PATTERNS AND PRODUCT SELECTION 52
15.1 WEAR PATTERNS FOR ABRASIVE AND IMPACT WEAR 52
15.2 WEAR PLATES AND GROUSER BARS 60
15.2.1 WEAR PLATES 60
15.2.2 GROUSER BARS (EG, BARS FOR REBUILDING WORN 60
16 SURFACING ALLOYS 62
16.1 CHROMIUM CARBIDE WEARFACING ALLOYS 62
16.2 WORK HARDENING ALLOYS (AUSTENITIC MANGANESE STEEL) 63
16.3 IRON BASED BUILD UP AND WEARFACING ALLOYS 64
16.4 TUNGSTEN CARBIDE WEARFACING ALLOYS 65
16.5 Ni BASED WEARFACING ALLOYS 66
17 GRADING OF WEAR RESISTANCE OF HARDFACING ALLOYS 67
18 METHODS OF WEAR PROTECTION - SUMMARY 68
19 REFERENCES 69

WELDING AND COATING METALLURGY
1 INTRODUCTION
Steels form the largest group of commercially important alloys for several reasons:
♦ The great abundance of iron in the earth’s crust
♦ The relative ease of extraction and low cost
♦ The wide range of properties that can be achieved as a result of solid state transformation such as
alloying and heat treatment
1.1 ROLE OF CARBON IN STEEL
Steels are alloys of iron with generally less than 1% carbon plus a wide range of other elements. Some of
these elements are added deliberately to impart special properties and others are impurities not completely
removed (sometimes deliberately) during the steel making process. Elements may be present in solid
solution or combined as intermetallic compounds with iron, carbon or other elements. Some elements,
namely carbon, nitrogen, boron and hydrogen, form interstitial solutions with iron whereas others such as
manganese and silicon form substitutional solutions. Beyond the limit of solubility these elements may
also form intermetallic compounds with iron or other elements. Carbon has a major role in a steels
mechanical properties and its intended use as illustrated in Figure 1.
As the carbon concentration is increases carbon steel, in general, becomes stronger, harder but less ductile.
This is an important factor when a steel is required to be welded by joining or surfacing.

Figure 1 Role of Carbon in Steel
1.2 WELDABILITY OF STEELS
When considering a weld, the engineer is concerned with many factors such as design, physical
properties, restraint, welding process, fitness-for-purpose etc., which can conveniently be summarized as
the base materials “weldability”. Weldability can be defined as “the capacity of a metal to be welded
under the fabrication conditions imposed into a specific, suitably designed structure, and to perform
satisfactorily in the intended service.”
Welding is one of the most important and versatile means of fabrication and joining available to industry.
Plain carbon steels, high strength low alloy (HSLA) steels, quench and tempered (Q&T) steels, stainless
steels, cast irons, as well as a great many non-ferrous alloys such as aluminium, nickel and copper are
welded extensively. Welding is of great economic importance, because it is one of the most important
tools available to engineers in his efforts to reduce production, fabrication and maintenance costs.
A sound knowledge of what is meant by the word “weld” is essential to an understanding of both welding
and weldability. A weld can be defined as a union between pieces of metal at faces rendered plastic or
liquid by heat, or pressure, or both, with or without the use of filler metal. Welds in which melting
occurs are the most common. The great majority of steels welded today consist of low to medium carbon

steel (less than 0.4%C).Practical experience over many years has proved that not all steels are welded with
ease. For example, low carbon steels of less than 0.15%C can be easily welded by nearly all welding
processes with generally high quality results. The welding of higher carbon steels or relatively thick
sections may or may not require extra precaution. The degree of precaution necessary to obtain good
quality welds in carbon and alloy steels varies considerably. The welding procedure has to take into
consideration various factors so that the welding operation has minimal affect on the mechanical
properties and microstructure of the base metal.
The application of heat, generally considered essential in a welding operation, produces a variety of
structural, thermal and mechanical effects on the base metal being welded and on the filler metal being
added in making the weld. Effects include:
♦ Expansion and contraction (thermal stresses etc.)
♦ Metallurgical changes (grain growth etc.)
♦ Compositional changes (diffusion effects etc.)
In the completed weld these effects may change the intended base metal characteristics such as strength,
ductility, notch toughness and corrosion resistance. Additionally, the completed weld may include defects
such as cracks, porosity, and inclusions in the base metal, heat affected zone (HAZ) and weld metal itself.
These effects of welding on any given steel are minimized or eliminated through changes in the detailed
welding techniques involved in producing the weld.
It is important to realize that the suitability of a repair weld on a component or structure for a specific
service condition depends upon several factors:
♦ Original design of the structure, including welded joints
♦ The properties and characteristics of the base metal near to and away from the intended welds
♦ The properties and characteristics of the weld material
♦ Post Weld Heat Treatment (PWHT) may not be possible
As discussed, a steels weldability will be dependent upon many factors but the amount of carbon will be a
principal factor. A steels weldability can be categorized by its carbon content as shown in Table 1.

Table 1 Common Names and their Typical Uses for Carbon Steel
COMMON NAME %C TYPICAL
HARDNESS
TYPICAL USE WELDABILITY
Low C steel <0.15 60 Rb Sheet, strip, plate Excellent
Mild Steel 0.15 – 0.30 90 Rb Structural shapes,
plate, bar
Good
Medium C Steel 0.30 – 0.50 25 Rc Machine parts, tools Fair
(preheat & postheat
normally required;
low H2
recommended)
High C Steel 0.50 – 1.00 40 Rc Springs, dies, rails Poor – Fair
(preheat and post
heat; low H2
recommended)
In order to understand the physical and chemical changes that occurs in steels when they are welded, a
basic understanding of the metallurgy of steels is necessary.
2 CRYSTAL STRUCTURE
Iron has the special property of existing in different crystallographic forms in the solid state. Below
910°C the structure is body-centred cubic (bcc). Between 910°C and 1390°C iron changes to a face-
centred cubic (fcc) structure.
Figure 2 Transformation of crystal structure for iron showing contraction occurring at 910°C.

Figure 3 BCC Crystal Structure
Figure 4 FCC Crystal Structure
Above 1390°C and up to the melting point at 1534°C the structure reverts back to body-centred
cubic form. These are known as allotropic forms of iron. The face-centred cubic form is a close-
packed structure being more dense than the body-centred cubic form. Consequently iron will
actually contract as it is heated above 910°C when the structure transformation takes place.
2.1 SOLUBILITY OF CARBON
The solubility of carbon in the bcc form of iron is very small, the maximum solubility being only
about 0.02 wt.% at 723°C. Figure 5 shows there is negligible solubility of carbon in iron at
ambient temperature (less than 0.0001 wt.%). Since steels nearly always have more carbon than
this, the excess carbon is not in solution but present as the intermetallic compound iron-carbide
Fe3C known as cementite.
`
WELDING AND COATING METALLURGY2 8 of 69

Figure 5 Solubility of carbon in α (bcc) iron as a function of temperature
In contrast the fcc form of iron dissolves up to 2% carbon, well in excess of the usual carbon
content of steels. A steel can therefore be heated to a temperature at which the structure changes
from bcc to fcc and all the carbon goes into solution. The way in which carbon is obliged to
redistribute itself upon cooling back below the transformation temperature is the origin of the wide
range of properties achievable in steels.
3 IRON - IRON CARBIDE PHASE DIAGRAM
Fundamental to a study of steel metallurgy is an understanding of the iron – iron carbide phase
diagram. The diagram commonly studied is actually the metastable iron – iron carbide system.
The true stable form of carbon is graphite, but except for cast irons this only occurs after
prolonged heating. Since the carbon in steels is normally present as iron carbide, it is this system
that is considered. Figure 7 shows the iron – iron carbide system up to 6 wt.% carbon. We will
now consider several important features of this diagram.

Figure 6 The iron-iron carbide equilibrium phase diagram
A eutectic is formed at 4.3% carbon. At 1147°C liquid of this composition will transform to two
solid phases (austenite + cementite) on cooling. This region is important when discussing cast
irons but is not relevant to steels.
TALLURGY2
3.1 AUSTENITE (γ)
This region in which iron is fcc, identified in Figures 7 and 8, dissolves up to 2% carbon. This
phase is termed austenite or gamma phase. With no carbon present it begins at 910°C on heating
but with 0.8% carbon it starts at 723°C. When a steel is heated into the austenite region all carbon
and most other compounds dissolve to form a single phase (i.e. normalizing).
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Figure 7 The austenite region of the iron-iron carbide diagram showing
maximum solubility of up to 2%C
3.2 FERRITE (α)
The region shown in Figure 9 where carbon is dissolved in bcc iron is very narrow, extending to
only 0.02% carbon at 723°C. This phase is termed ferrite or alpha phase. Although the carbon
content of ferrite is very low other elements may dissolve appreciably in it so ferrite cannot be
considered as “pure iron”.
Figure 8 The ferrite region of the iron-iron carbide diagram
3.3 PERITECTIC
The region at the top left portion of the phase diagram enlarged in Figure 10 is where the iron
reverts back to the bcc structure known as delta ferrite. Here again the solubility for carbon is
low, only 0.1 wt.% at 1493°C. The part of the diagram at 0.16% carbon having the appearance of
an inverted eutectoid is called a peritectic. At this point a two phase mixture of liquid and solid
(austenite) transforms on cooling to a single phase solid of austenite. This portion of the phase
diagram will not be discussed in detail, but it should be recognized since it has been invoked to
explain various hot cracking phenomena in welding.

Figure 9 Peretectic region of the iron-iron carbide diagram
3.4 PEARLITE
At 0.8% carbon and 723°C a eutectoid is formed as illustrated in Figure 11. This is similar to the
eutectic transformation but involves a solid phase transforming into two different phases on
cooling (ferrite and cementite). This eutectoid mixture is called pearlite. Figure 12 shows how
the two phase constituents that make up pearlite are formed. Note that pearlite is only one of
many phases that can be produced from ferrite and cementite (depending on cooling rate).
Cementite (iron carbide) itself is very hard - about 1150 Hv – but when mixed with the soft ferrite
layers to form pearlite, the average hardness of pearlite is considerably less.
Figure 10 The eutectoid point on the iron-iron carbide diagram

Figure 11 Schematic View of how pearlite is formed in an approx. 0.4%C steel
This region of the phase diagram (where carbon concentration is less than 0.8%) is of the most
interest to a study of steels and their weldability which will be discussed in more detail later.
3.4.1 PEARLITE GROWTH
A steel with 0.8 wt.% carbon, it will be recalled, transforms on cooling through 723°C to the two
phase eutectoid constituent pearlite. In pearlite the two phases ferrite and cementite are mixed
closely together in fine layers. As the ferrite contains very little carbon while the cementite has
6.7%, carbon atoms must diffuse to the growing cementite plates as shown in Figure 13.
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Figure 12 Schematic View of different pearlite growth rates
The distance they can diffuse, and hence the spacing of the plates, depends on how fast the pearlite
is growing. A fast growth rate means less time for diffusion and a finer pearlite results. Figure 14
shows a typical pearlite microstructure.
Figure 13 Typical Lamellar Appearance of Pearlite. Mag:X1500
3.5 PRO-EUTECTOID FERRITE
If the steel has less than 0.8 wt.% carbon (termed hypo-eutectoid steel) ferrite will be formed first
from the austenite. The example in Figure 15 shows a steel of 0.4 wt.% carbon. This ferrite is
called pro-eutectoid ferrite because it transforms first on cooling as illustrated in Figure 15. As
transformation continues and the temperature drops, the remaining austenite becomes richer in
carbon. At 723°C the steel comprises ferrite and the remaining austenite (which contains 0.8wt.%
carbon). With further cooling, the austenite then transforms to pearlite producing a final structure
in the steel of pro-eutectoid ferrite and pearlite.
Figure 14 Phase Transformation on Cooling a 0.4%C Steel

The amounts of pro-eutectoid ferrite and pearlite can be estimated by application of the lever rule
(see references for more detailed information). For a 0.4 wt.% carbon steel about 50% will be
ferrite and 50% pearlite. Similarly a steel of more than 0.8 wt.% carbon (from 0.8 wt.% up to 1.8
wt.% carbon is termed hyper-eutectoid steel) first transforms to cementite (i.e. pro-eutectoid
carbide) with the remaining austenite forming pearlite as shown in Figure 16.
Figure 15 Phase Transformation on Cooling a 1.2%C Steel
3.6 PHASE TRANSFORMATIONS IN LOW ALLOY STEELS
Figure 17 shows the appearance of a polished and etched section of an approximately 0.6wt.%
carbon steel. You can see that the pro-eutectoid ferrite has formed initially at the austenite grain
boundaries, nucleation taking place at several points around each austenite grain. Since each
region of ferrite becomes an individual grain, its grain size will be very much smaller than that of
the parent austenite. Ferrite continues to form and grow until the final transformation of
remaining austenite to pearlite. The ferrite does not always appear as neat, equiaxed grains as
shown in Figure 17, but can occur as long spikes from the grain boundaries or even nucleate
within the austenite grain. This can occur quite markedly from the welding process due to the
cooling rates imposed by the heat input (i.e. travel speed).

Figure 16 Prior Austenite Boundaries Showing Pro-Eutectoid Ferrite
On reheating the steel the process reverses and the pearlite and ferrite grains transform back into
single phase austenite to form completely new grains. The temperature required to get complete
transformation depends on the carbon level as seen from the phase diagram (see Figures 7 and
15)and ranges from 910°C for zero carbon to 723°C for 0.8 wt.% carbon.
3.7 GRAIN GROWTH
Heating to higher temperatures than those necessary to get complete transformation causes the
austenite grains to grow. The final size of the austenite grains depends not only on the
temperature reached but also on the type of steel. Some steels containing small precipitates such
as aluminium and vanadium nitride retain small grain size up to high temperatures. These are
known as fine grained steels. Steels can be deliberately made as coarse grain or fine grain. Fine
grained steels are tougher and are more commonly specified for most structural applications.
The effect of austenizing temperature on grain size is shown in Figure 18. It shows that although
grain growth is restricted in a fine grain steel, at a sufficiently high temperature the precipitates
dissolve and the steel behaves as a coarse grain steel. Thus at sufficiently high temperature, grain
growth can occur with subsequent loss of toughness. This is an important consideration in the
HAZ associated with welding.

Figure 17 Schematic Effect of Temperature on Grain Growth for Coarse and Fine Grained Steels
3.8 NON-EQUILIBRIUM COOLING
The phases and microstructures predicted by the iron – iron carbide diagram occur in steels cooled
very slowly. In addition the diagram assumes that carbon is the only alloying element present in
the steel. With the addition of other common alloying elements such as manganese, silicon,
nickel, titanium, molybdenum, chromium etc., the phase diagram can still be used except that it
will be distorted and the lines may move to slightly different locations.
Figure 18 Effect of Various Element Additions on the Recrystallization Temperature
For example the presence of alloy elements changes the recrystallization (eutectoid) temperature
as shown in Figure 19. In structural steels the concentration of alloys is generally quite small
(austenitic manganese steels are an exception containing over 12 wt.% manganese) and the basic
iron – iron carbide phase diagram is not distorted very much from equilibrium conditions.

3.9 MARTENSITE - EFFECT OF RAPID COOLING
The rate of cooling has a major effect on the types of microstructures formed and unless the steel
cools slowly the iron – iron carbide phase diagram cannot be used. The reason is that the
transformation of austenite to pearlite requires the diffusion of carbon to the sites of growing
carbon, a process which takes time. We saw how a faster cooling rate produced finer pearlite.
With even faster cooling rates less time is available for diffusion and pearlite cannot form.
Alternative microstructures form with their exact morphology depending on just how quickly the
steel cools. In a water quench, for example, the cooling rate is so rapid there is no time for any
diffusion, and the carbon remains trapped in the same place as it was in the austenite. A rapid
quench cannot suppress the crystal structure change from fcc to bcc but the presence of trapped
carbon in the bcc phase distorts it to a tetragonal shape, as indicated in Figure 20, rather than a true
cubic structure. This is called martensite.
Figure 19 Schematic Transformation of Austenite (BCC) To Martensite (Tetragonal) With Increasing %C
The amount of carbon influences the amount of distortion in the crystal structure as shown in
Figure 20. This in turn affects the hardness of the martensite as shown in Figure 22.
Under the microscope as shown in Figure 21 martensite has the appearance of a mass of needles.

Figure 20 Martensite Microstructure
Martensite can be very hard and brittle when it contains appreciable amounts of carbon. The
hardness
depends almost exclusively on the carbon
content with other elements having little effect as illustrated in
Figure 22.
Figure 21 Effect of Carbon and Alloying on the Hardness of Martensite
The formation of martensite can occur in the HAZ adjacent to a weld deposit due to the fast
cooling rates imposed by the welding process. This is discussed in more detail in Section !!
3.10 BAINITE
Intermediate between a rapid quench that produces martensite, and a slow cool producing pearlite,
other constituents may form particularly in alloy steels. The most important of these is bainite.
Figure 22 Microstructure of lower bainite (X1000)

Bainite is still a two phase mixture of ferrite and iron carbide but unlike the cementite plates in
pearlite the carbide in bainite is spherical. Bainite formed above 300°C contains relatively coarse
particles of the Fe3C form of iron carbide (cementite) and is termed upper bainite. When formed
below 300°C bainite has a much finer structure with the carbides tending to form striations across
the ferrite laths. This is termed lower bainite. The carbides in lower bainite are Fe2.4C known as
epsilon (ε) carbide. Some steels in the bainitic condition may possess ductility and toughness
superior to that shown by the same steel in the Q&T condition.
4 TRANSFORMATION DIAGRAMS
Since the iron – iron carbide phase diagram is only valid for very slow cooling rates, alternative
diagrams for determining the constituents present in a more rapidly cooled steel have been
developed. There are two types:
♦ Time Temperature Transformation (TTT) curves where the steel sample is held at a constant
temperature until transformation is complete.
♦ Continuous Cooling Transformation (CCT) curves where the steel sample is cooled from the
austenitic region at different cooling rates.
Although these diagrams are principally designed for the foundry metallurgist and heat treater etc.,
they are an excellent tool for use by welding engineers where fast cooling rates need to be
evaluated near to the welded area.
4.1 TIME TEMPERATURE TRANSFORMATION (TTT) DIAGRAMS
Consider heating a sample of steel until it is fully austenitic then quenched to some temperature
below the equilibrium transformation temperature as shown in Figure 24.
Figure 23 Schematic Representation of TTT
If we hold the steel at this temperature we find there is a delay before transformation begins and a
further elapse of time while transformation takes place. The delay depends on the temperature at

which the steel is held and we can plot this information on a diagram of temperature against time
for a given steel composition.
Figure 24 Schematic TTT Curve for Carbon Steel
An example of such a time-temperature-transformation (TTT) diagram for a carbon steel is shown
in Figure 25. Note that at high temperatures (Figure 26) the steel transforms to pro-eutectoid
ferrite followed by pearlite.
Figure 25 TTT Curve Illustrating High Temperature Transformation of Pro-Eutectoid Ferrite

At lower temperatures less pronounced pro-eutectoid ferrite is formed and the pearlite is finer. At
about 550°C the pearlite forms in the shortest time and there is no pro-eutectoid ferrite (Figure
27).
Figure 26 TTT Curve Illustrating Pearlite Transformation
Cooling down to below this range (approximately 450°C) transformation to bainite occurs, taking
a longer time for lower temperatures (Figure 28).
Figure 27 TTT Curve Illustrating Transformation to Bainite
At a sufficiently fast cooling down to low temperature martensite can begin to form (Figure 29).
Note that it forms almost instantaneously and does not grow as a function of time. For each steel
specification there is a fixed temperature Ms at which martensite starts to form and a fixed
temperature Mf at which transformation is complete. The percentage of martensite formed
therefore depends only on the temperature to which the steel is rapidly cooled c and not on how
long it is held there. If the composition of the steel is known, the Ms temperature can be calculated

(see Section !!!). Note that for some compositions the Mf temperature can be below ambient
temperature.
Figure 28 TTT Curve Illustrating Martensite Formation
4.2 CONTINUOUS COOLING TRANSFORMATION (CCT) DIAGRAMS
Now consider the case of continuous cooling. We may superimpose a cooling curve on the TTT
diagram as illustrated in Figure 30 in order to get an idea of what microstructures form, but it is
more accurate to use a diagram established under continuous cooling conditions. The CCT
diagram is slightly different from the TTT curve.
Figure 29 Cooling Curves Superimposed onto TTT Curve for Typical Carbon Steel

4.2.1 CRITICAL COOLING RATES
You should note that in plain carbon steels bainite generally will not form during continuous
cooling because of the shape of the TTT diagram. The bainite region is tucked under the pearlite
area so a cooling curve either hits the pearlite curve or misses it completely as shown in Figure 30.
At cooling rates fast enough to miss the nose of the curve martensite is formed.
This is an important concept since the cooling rate at which martensite can form in a HAZ
strongly influences the risk of cracking during welding and gives an indication of a steels
“weldability”. This will be discussed in more detail in Section 8.
4.2.2 DETERMINING CCT DIAGRAMS
The exact shape of a CCT curve depends on the chemistry of the steel and on the heating and
cooling cycles. CCT diagrams are available for numerous carbon and alloy steels and if desired
can even be established for specific weld metal.
4.3 EFFECT OF ALLOYING ELEMENTS
Alloy elements have significant effects on the shape of the CCT and TTT diagrams which allow
different microstructures to be produced in alloy steels. Chromium and molybdenum, for
example, shift the top (pearlite) part of the curve to the right i.e. to longer times, thus exposing the
bainite region. Steels containing these elements such as 4135 can produce bainite on continuous
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cooling.
Figure 30 CCT Curve for 4135 Steel
The entire TTT curve may also shift to the right with additions of certain elements (e.g.
chromium, vanadium, molybdenum and others) to greater times allowing martensite to form at
much slower cooling rates. This increases the “hardenability” of the steel, but also increases the
risk of cracking from welding if proper precautions are not taken.
4.4 MS AND MF TEMPERATURES
The other notable effect of alloy element addition is to change the martensite start (Ms) and
martensite finish (Mf) temperatures. Increasing the carbon content, for example, depresses the Ms
to lower temperatures as shown in Figure 32.
Figure 31 Schematic Diagram Showing the Influence of %C on Martensitic Start Temperature
Other elements affect martensite formation and the combined affect can be approximated by the
following equation:
Ms (°C) = 550 – 350 ×%C - 40×%Mn - 35×%V - 20×%Cr - 17×%Ni -10×%Cu - 10×%Mo - 5×%W + 15×%Co +
30×%Al
As mentioned earlier, if the Mf is below ambient temperature, martensite transformation is not
completed and the steel contains retained austenite.
Note that since the product of various alloy element additions affects the Ms, we can affect a steels
hardenability by using small quantities of several alloy elements rather than a large quantity of one
element, for example, carbon. This is important when designing a steel for not only hardenability
but also its weldability.

5 HARDENABILITY / WELDABILITY OF STEELS
Figure 32 Correlation of CCT and TTT Diagrams With Jominy Hardenability Test Data
for an 8630 Type Steel
The hardenability of steels can be determined by performing a Jominy end quench test. The alloy
steel test specimen is a cylinder one inch diameter and four inches long, which is heated to the
austenitic region (above 910°C) then placed in a fixture where it is quenched by water or brine
impinging on one end. The fastest cooling rate occurs at the bar surface in contact with the water
jet with progressively slower cooling rates being experienced away from the end. Thus the
microstructure formed in the surface region could be martensitic with high hardness and the
interior could be pearlitic with no hardening at all. The depth to which a steel hardens is a measure
of its hardenability. If we add alloying elements that allows deeper hardening, then that steel is
said to have higher hardenability. This is important, for example, when considering mechanical
properties and weldability of such a steel. Hardness tests are commonly used on Jominy samples
to determine that steels hardenability.
Figure 33 illustrates the TTT diagram for a common chrome-molybdenum steel (4137) with a
Jominy end quench test superimposed. Thus the microstructure and hardness can be correlated on
the one diagram.
The cooling rate curves represent the same cooling rate conditions located along the Jominy end-
quench test bar. At the top of Figure 33, the measured hardness curve has been superimposed over

a schematic of the end-quenched bar. Four representative locations (A, B, C, D) along the bar have
been related to the representative cooling curves(CCT) and isothermal transformation (TTT)
curves. Thus location A on the bar experienced a fast cooling rate resulting in austenite
transforming to martensite producing the high hardness indicated. Similar cooling rate effects
need to be considered from a weldability viewpoint.
The addition of alloying elements (for example Mo, Cr, Mn) to steel increases the hardenability by
slowing down the rate of austenite transformation. The data is plotted as shown in Figure 34 for a
0.45%C steel with different alloying additions.
Figure 33 Typical End-Quench Curves for Several 0.45%C Low Alloy Steels
Several formulae have been developed which assign a contributing factor to each element addition
and its effect on hardenability and conversely weldability. The maximum hardness attainable (and
therefore its weldability characteristics) in carbon and low-alloy steels, however, is still almost
exclusively dependent upon the carbon content.

5.1 CARBON EQUIVALENT (CE) & WELDABILITY
Depth of hardening is not a relevant concept in a welding situation, but we are interested in the
hardness produced at a given cooling rate or the critical cooling rate to produce a given hardness
in the HAZ of a weld. There are several models that have been developed to calculate
hardenability from a welding process. The simplest model is one in which the effects of individual
alloying elements are added together (a linear model) to produce a carbon equivalent (CE) which
in turn relates to a critical cooling rate to produce a given hardness. Figure 35 shows a reasonable
correlation between the CE plotted against critical cooling rate from 540°C to give a hardness of
350Hv in the HAZ.
Figure 34 Linear Correlation of CE and Cooling Rate for a Fixed HAZ Hardness
Another linear model has been used to predict the hardness of the HAZ for different cooling rates
in low alloy steels and is illustrated in Figure 36.
COATING METALLURGY2 28 of 69
0
200
400
600
800
1000
1200
1400
0.62 0.72 0.82 0.92 1.02
CE
Hv
Hv@50deg/sec
Hv@100deg/sec
Hv@200deg/sec
Hv@500deg/sec
WELDING AND

Figure 35 Correlation of HAZ Hardness and CE as a Function of Cooling Rate
CE’s are used widely in industry as measures of weldability. Several different formulae have been
developed and some are even incorporated into national codes and specifications. In general terms,
other factors being equal, as the carbon content increases, so does the difficulty in weldability. In
practice, this means generally using higher preheats until cracking and restraint problems are
overcome.
Using an engineering/analytical approach becomes very useful when confronted with unknown
material compositions, and weld repairs can become challenging where reverse engineering must
be utilized to develop a repair procedure. The engineering approach may involve evaluating
composition, hardenability, service conditions, size, restraint conditions, and PWHT feasibility.
One of the popular methods for determining weldability is to review the hardenability of the base
material. As discussed earlier the CE formula(s) have been developed as a convenient method of
normalizing the chemical composition of a material into a single number to indicate its
hardenability. Review of the literature indicates no less than a dozen different formulas have been
developed. One of the most commonly used formulas for calculating the CE is the IIW formula
(shown in Figure 36):
15
Cu
Ni
5
V
Mo
Cr
6
Mn
C
CE
+
+
+
+
+
+
=
It must be stated that low carbon steel and carbon – manganese steels generally behave in a
predictable manner and are successfully welded with preheat and PWHT criteria outlined in codes
such as
AWS D1.1, Structural Welding Code – Steel. The CE is not usually evaluated on these materials.
Medium carbon, HSLA, and Q&T Steels, however, present different challenges where
consideration of CE, restraint, hydrogen control, PWHT not practicable, weld filler chemistry
mismatch, weld heat input etc. can be critical to successful repair welding. These factors can be
summed up as a materials weldability, and it is these factors that will be considered in Section 8.
5.2 TEMPERING – EFFECTS OF REHEATING
As discussed earlier martensite produced in a quenched steel is hard and brittle and in most cases
the steel is unusable in that form. The toughness may be improved by a process of tempering. This
involves reheating the steel to below the transformation temperature (723°C), holding for a period
of time, then cooling to ambient temperature as illustrated in Figure 37. During tempering the
carbon trapped as an interstitial in the martensitic tetragonal structure is released. Carbon atoms
diffuse and precipitate as small carbides. With enough time and at sufficiently high temperatures
cementite (Fe3C) forms, not as plates as in pearlite, but as spherical particles. This microstructure
is known as bainite(see section 3.9).

Figure 36 TTT Curve illustrating Q&T process to form Bainite
Improvement in toughness is accompanied by a loss of hardness which is a function of both
temperature and time (however temperature is more effective – the higher the temperature the
faster the tempering transformation as illustrated in Figure 37). The temperatures typically
selected for post weld heat treating or stress relieving welded steel are generally high enough to
cause rapid tempering of the HAZ.
5.3 SECONDARY HARDENING
In some steels containing specific alloy elements tempering may actually cause an increase in
hardness as the tempering temperature is raised as shown in Figure 38. This is known as
secondary hardening and is caused by strong carbide forming elements such as molybdenum,
chromium, and tungsten combining with carbon to form alloy carbide precipitates in certain
temperature ranges. This behavior of secondary hardening is put to good use in the tempering of
tool steels such as high speed tool steels. When considering a weld repair on such steels, the
preheat and interpass temperatures is normally selected at a temperature below the secondary(or
tempering) temperature, particularly if PWHT is not practical.
Figure 37 Alloying Effect on Secondary Hardening

6 HYDROGEN CRACKING RELATED TO WELDABILITY
Hydrogen can embrittle a steel at both elevated and ambient temperatures. The term hot cracking
is used to signify that cracking has occurred at elevated temperature while cold cracking is used to
generally signify cracking in low alloy steel at ambient temperature. It was during World War 2
that it was realized that hydrogen dissolved in weld metal was one of the causes of cold cracking
in low alloy steel welded joints (i.e. the catastrophic failure of the welded Liberty ships). These
failures led to the development of low hydrogen electrodes which made possible successful
welding of the alloy steels used today.
Hydrogen pickup is derived from hydrogen containing chemical compounds that are dissociated in
the arc column. They can originate, for example, from contamination on the workpiece or from
moisture in the welding flux. It is the hydrogen sourced from electrode coatings or fluxes which is
the most important. Electrode coatings consist of minerals, organic matter, ferro-alloys, and iron
powder bonded with, for example, bentonite (a clay) and sodium silicate. The electrodes are baked
after coating, and the higher the baking temperature the lower the final moisture content of the
coating. Some electrode coatings may pick up moisture if exposed at ambient conditions (basic
coated electrodes). Where hydrogen cracking is a risk, special flux coatings are used to maintain
low hydrogen content. In practice, welding specifications stipulate the allowable moisture
content. It is, however, important to note that the method or welding procedure adopted as well as
the type of electrode flux used can affect the hydrogen content in a weld or HAZ.
With hot cracking, embrittlement occurs in carbon and low alloy steels by a chemical reaction
occurring between hydrogen and carbides which causes irreversible damage – either
decarburization or cracking or both. Of much greater importance in welding is hydrogen
entrapped in the weld or HAZ causing embrittlement. Hydrogen cracking can subsequently occur
at some later time (sometimes days) once a weld repair is complete, generally at service
temperatures between – 100°C and 200°C. This embrittlement is due to physical interactions
between hydrogen and the crystal lattice structure of the steel and is reversible by removal of
hydrogen by stress relieving allowing the ductility of the steel to revert back to normal. Hydrogen
cracking can occur in either the weld metal, HAZ, or base metal and be either transverse or
longitudinal to the weld axis. The level of preheat or other precautions necessary to avoid cracking
will depend on which region is the more sensitive. In carbon - manganese medium strength steels
the HAZ is usually the more critical region and weld metal rarely causes a problem.
Cracking due to dissolved hydrogen is now thought to occur by decohesion. Where there is a
defect, discontinuity or pre - existing crack and a tensile stress applied, hydrogen is considered to
diffuse preferentially to the region of greatest strain i.e. near to the stress concentration such as
near a crack tip. The presence of a relatively large concentration of hydrogen reduces the cohesive
energy of the crystal lattice structure to the extent that fracture occurs at or near the stress
concentrator. This view is consistent with observations that cracking can occur slowly (the crack
velocity being dependent on the diffusion rate of hydrogen) and is quite often discontinuous.
In welding, the region most susceptible to hydrogen cracking is that which is hardened to the
highest degree (areas where the welding residual stresses is greatest) although regions of coarse
grain growth can be a contributing factor. The most crack-sensitive microstructure is high carbon
martensite.

Hot or cold cracking in the weld metal or HAZ depends on the same fundamental factors as in the
base metal, i.e. hydrogen content, microstructure and residual stress. In practice the controlling
variables are usually strength, hydrogen content, restraint, stress concentrations, and heat input.
Figure 38 Minimizing Heat Input by Multi-Pass Welding
In single pass welds and root runs of multiple pass welds the root pass may provide a stress
concentration which can lead to longitudinal cracks in the weld metal. High dilution of the root
run (high heat input) can often result in a harder weld bead more likely to crack (this is commonly
seen in such applications as pipeline welding). Figure 40 illustrates the physical appearance of
hydrogen cracking in welds.
Figure 39 Schematic View of Typical Weld Cracks
Figure 40 Cracking Caused By Lack Of Fusion in Weld
In Figure 41 the crack has initiated at the root of the weld where a lack of fusion can be seen. The
crack has then traveled through the HAZ mainly in the coarse grained region. In heavy multiple –
pass welds cracking will generally be transverse to the weld direction, sometimes running through
the weld itself since the maximum cooling rate is along the weld axis. Many HSLA steels in
critical repair situations where PWHT is impracticable are welded using a filler metal of good
toughness and ductility and in such cases the HAZ may be more crack sensitive.
The risk of hydrogen-induced cold cracking in the weld can be minimized by:
♦ Reducing hydrogen pick-up (low hydrogen flux chemistries)
♦ Maintaining a low carbon content
♦ Avoiding excessive restraint
♦ Control of welding procedures (preheat; heat input; PWHT etc.)
♦ Developing a non-sensitive weld microstructure

In carbon or carbon-manganese steels (i.e., those with steep hardening curves as shown in Figure
35!!) welding conditions can be selected to avoid the cooling rates at which martensite is
produced. This could include preheat; high heat input welding; slow cooling etc.
In low alloy steels or those where a hard HAZ cannot be avoided, other steps must be taken to
prevent cracks. These often involve applying preheat and interpass temperatures to allow the
diffusion of hydrogen out of the weld metal. Figure 43 shows that quite moderate temperatures are
highly effective in removing hydrogen.
Figure 41 Effect of Moderate Postheat on Hydrogen Content in a Cooled Weld
The freedom of selecting a suitable welding solution is sometimes limited. The solution must be
practicable and economic. Further constraints may be applied by the job such as base metal
condition, size, location, PWHT not practicable, equipment availability etc. In such cases, the
welding engineer may need to consider the steels CE and Ms temperature by referring to its TTT
and CCT curves in providing a weld procedure.
6.1 LAMELLAR TEARING
Lamellar tearing is a form of cracking that occurs in the base metal of a weldment due to the
combination of high localized stress and low ductility of the base metal. It is associated with
regions under severe restraint, for example, tee and corner joints; heavy sections etc.
Figure 42 Example of Lamellar Tearing from Welding

The cracks appear close to or a few millimetres away from the HAZ at right angles to the weld
interface as shown in Figure 44. In HSLA steels that form martensite in the HAZ, hydrogen -
induced cold cracking will generally form preferentially, but in plain carbon steels of low
hardenability, hydrogen increases the susceptibility to lamellar tearing quite markedly due to HAZ
stresses. There is no correlation between heat input and the incidence of lamellar tearing, but in
the presence of hydrogen a low heat input might tip the balance towards hydrogen cracking
because of a lack of time for hydrogen to dissipate away from the weld area.
Lamellar tearing may, in principle, be avoided by:
♦ Design modification
♦ Buttering weld runs and temper bead welding
♦ Control of welding procedures (preheat; heat input; PWHT etc.)

7 REHEAT CRACKING IN THE HAZ
During PWHT, stress relief treatment, or during service at elevated temperature ( generally higher
than 400°C) reheat cracking may occur in the HAZ of welds in alloy steel.
Cracking is due to the combined effects of embrittlement and strain. At elevated temperature,
precipitation occurs within the HAZ grains but not at the grain boundaries, where there is a
denuded zone. Consequently, the interior of the grain is relatively hard and the boundary region
relatively soft. When the residual welding strain relaxes, the deformation is concentrated at the
weld boundaries. If the degree of strain exceeds the ductility of the grain boundary regions,
cracking will take place.
Cracks are intergranular and follow the prior austenite boundaries. They may start at high stress
points such as the toe of welds, an unfused root run, or may be sub-surface. Cracking can also
occur in highly restrained joints such as in very heavy components.
Factors that contribute to reheat cracking are:
♦ Susceptible alloy chemistry
♦ Susceptible HAZ microstructure
♦ High level of triaxial residual strain
♦ Temperature in the strain relaxation (creep) range
Most alloy steels suffer some degree of embrittlement in the coarse-grained region of the HAZ
when heated at elevated temperatures (e.g. 600°C). Elements that promote such embrittlement are
Cr, Cu, Mo, B, V, Nb, and Ti. Molybdenum-vanadium and molybdenum-boron steels are
particularly susceptible especially if the vanadium content is over 0.1%. The relative effect of the
various elements has been expressed quantitatively in one formula due to Ito:
Psr = Cr + Cu + 2Mo + 10V + 7Nb + Ti – 2
When Psr is ≥ 0 stress relief or reheat cracks may occur.
This formula does not work for low carbon(<0.1%C) or high chromium (>1.5%Cr), these steels
being resistant to reheat cracking. The use of low heat-input processes (MMAW or GMAW) and a
weld metal of high yield strength and a high degree of toughness are important benefits.
Reheat cracks may also form:
♦ When welding dissimilar steels due to differential thermal expansion coefficients

♦ At elevated service temperature and high stress loading
Reheat cracking can be avoided by:
♦ Proper base metal selection (using the Ito formula)
♦ Designing to minimize restraint (eliminate stress concentrators etc.)
♦ Using low heat input weld process
♦ Use higher preheat/interpass temperature
♦ PWHT after part-welded

8 WELDING STEELS CONSIDERED DIFFICULT
Welding of HSLA and Q&T steels may pose several problems and a careful study of the steel and
its intended application is necessary before specifying a welding procedure. Sometimes welding is
required after the steel has been heat treated i.e. Q&T making it almost impossible to achieve
uniform properties across the welded joint. In such cases it is preferable to carry out the welding
prior to the Q&T operation. The weld metal in this case must be selected to have matching
chemical properties so that as near as possible, uniform properties are achieved after the Q&T heat
treatment. To the maintenance engineer this is sometimes not possible to carry out for several
reasons such as size, in situ, economics etc. In these cases, the weld metal and repair procedure
have to be carefully considered to ensure, as near as possible, matching mechanical properties are
obtained (or superior) with HAZ hardness taken into account.
8.1 PROCEDURAL CONSIDERATIONS
To prevent martensite from forming during welding, sufficient preheat must be applied to the
component to hold it above the Ms temperature until welding is complete. All deposited weld
metal and the HAZ remain austenitic during the welding operation and transform together on
cooling to produce a uniform structure. In applying this approach the TTT diagram for the steel
can be studied to determine the preheat temperature, the maximum allowable time for completion
of welding, and the cooling rate required. The preheat and interpass temperatures thus selected
must also be below the tempering temperature of the base metal in order to maintain its
mechanical properties. The weld metal selected must provide adequate strength and toughness
and, if necessary, without the benefit of a subsequent PWHT.
8.2 POST WELD HEAT TREATMENT (PWHT)
It is common practice to apply a PWHT or stress relief to temper the welded joint and soften the
HAZ. Additionally PWHT removes hydrogen and lowers residual stresses imposed by the service
conditions and the welding operation. In the case of a Q&T steel the PWHT must not be higher
than the original tempering temperature otherwise a loss of physical properties such as strength
could occur dropping it below specification. To reduce the risk of cracking the PWHT may be
carried out immediately after welding is completed without letting the component cool down or
carried out several times during the weld repair which can be costly. As discussed previously, it
may be impractical to carry out PWHT. The weld repair procedure needs to be carefully
considered to minimize the possibility of cracking in service by ensuring that the welded
component has “fitness-for-purpose”.
8.3 THE HEAT AFFECTED ZONE (HAZ)

The HAZ undergoes a complete thermal cycle which determines the microstructure. Grain growth
is an important factor in the HAZ and the weld. In the HAZ of a coarse grained steel there is a
wide region where grain growth has occurred but in a fine grained steel, grain growth is resisted
except in the narrow region immediately adjacent to the weld fusion boundary where temperatures
are very high. Fig.7.10 shows an example of grain growth in a welded joint.
75mm
Figure 43 Macrosection of High Input Weld Showing Coarse Grain Size
The type of microstructure formed in the coarse-grained region of a steel depends upon:
♦ The carbon content
♦ The alloy content
♦ The time at elevated temperature
♦ The cooling rate
For any given steel, the greater the weld heat input the longer the time spent above the grain
coarsening temperature of the steel, and the coarser the grain size. Steels containing grain refining
additions such as titanium, vanadium, niobium, and aluminium are exceptions in that a fine HAZ
grain size may be achieved right up to the fusion boundary. Titanium nitride is very stable and
may not completely dissolve in the HAZ even at the temperatures immediately adjacent to the
fusion boundary. This can be advantageous with high heat input welds such as submerged arc
welding.
Figure 44 illustrates four welds in a carbon steel that have been welded with different heat inputs.

Figure 44 Correlation of Heat Input and HAZ Hardness
Alongside each weld the HAZ transforms to a microstructure dependent on the cooling rate of that
weld. For higher heat input welds, the cooling rate will be slower. In Figure !!, for the small
rapidly cooled welds, martensite is formed. For the large, slowly cooled welds the HAZ structure
is pearlite. The hardness of the HAZ is much higher in those welds in which martensite is present
as illustrated.
Adjacent to the weld the base metal undergoes various changes according to peak temperature and
cooling rate experienced at various locations away from the weld joint. Close to the fusion zone
the peak temperature will be high enough to cause complete transformation to austenite and some
grain growth.
Figure 45 Effect of Welding on Grain growth Relative to the Iron – Iron Carbide Phase Diagram
At some distance away from the fusion zone the temperature is not sufficient to cause any
microstructural changes although other effects such as strain aging (plastic deformation) may
occur. In between a range of mixed structures may be observed as illustrated in Fig.38.

The austenite grain size in the HAZ is controlled mainly by the weld heat input1
but it is also influenced by the
shape of the fusion zone.
8.3.1 LOSS OF TOUGHNESS IN THE HAZ
The microstructure itself may have an effect on the crack sensitivity or toughness in the HAZ.
Certain factors are known to lower the toughness of the HAZ:
♦ Grain size – An increasing austenite grain size in the HAZ is likely to result in lower
toughness. Grain size is determined largely by heat input and base metal chemistry.
♦ Heat Input – An increasing heat input caused by welding amperage; arc process;
weaving etc. can result in lower toughness. Indeed some high strength low alloy
(HSLA) steels specify heat input requirements for weld joining.
♦ Precipitation Hardening – From the presence of micro-alloy elements. Again
precipitation is encouraged by high heat input because of the longer times at high
temperatures and the slower cooling rates.
♦ Plastic Deformation – The contraction of a cooling weld may cause plastic deformation
in certain parts of the HAZ, particularly around any residing defects (such as nitrogen,
sulfides etc.), with consequent loss of toughness.
♦ Post weld heat treatment (stress relief) of micro-alloyed steel can cause a considerable
amount of precipitation of fine carbides with a substantial decrease in toughness in the
HAZ.
In practical terms, restrictions on heat input may mean some welding processes such as
electroslag, submerged arc, and flux-cored arc cannot be used. Other restrictions such as preheat;
interpass temperature; and width of weave would need to be considered also.
8.4 PREHEAT & CARBON EQUIVALENT
The preheat temperature required depends on the susceptibility of the HAZ to hydrogen cracking,
and much research has been done to find compositional formulae to indicate this. One formula for
calculating the preheat for welding of structural low alloy steels is given below:
15
Cu
Ni
5
V
Mo
Cr
6
Mn
C
CE
+
+
+
+
+
+
=
sec
/
mm
A
V
HeatInput
1000
×
×
=
1

( ) 0.25)
(CE
350
C
P o
TEMP −
=
PREHEAT TEMPERATURE FOR LOW ALLOY
STEELS ONLY
0
50
100
150
200
250
300
350
0
.
2
5
0
.
3
0
.
3
5
0
.
4
0
.
4
5
0
.
5
0
.
5
5
0
.
6
0
.
6
5
0
.
7
0
.
7
5
0
.
8
0
.
8
5
0
.
9
0
.
9
5
1
CE
P(TEMP)
Figure 46 Preheat Temperature as a Function of CE
8.4.1 SEFERIAN GRAPH
The Seferian graph shown in Fig. 43 takes into account CE, and restraint in calculating preheat.

9 SUMMARY
Modern structural steels with their demands for strength, toughness, and good welding behavior
have evolved to depend less on carbon content as a strengthening agent and more on fine grain
size and precipitation hardening. This has meant that welding (specifically weld repair) procedures
may now have to utilize consumables which meet stringent property requirements as well as
avoiding cracking and other defects during welding.

10 GENERALASPECTS CONCERNED WITH WEAR
PROTECTIVE COATINGS
There are a series of questions that must be asked when approaching the subject of wear surfacing.
The main ones being:
♦ Cause and Type of Wear?
♦ Base Metal?
♦ Area/Thickness?
♦ Facilities Available?
♦ Preheat? (Depends upon the base metal and its thickness)
These are perhaps the five most obvious questions that must be asked and answered. There is also
the question of suiting the wear resistant alloy to practical aspects and to give optimum economic
reward.
For example if the wearing parts of a machine are replaced only every month it is unsatisfactory to
increase wearing life to six weeks as this does not correspond to the shutdown time. In this case it
will be necessary to select an alloy that doubles the working life as a minimum requirement to the
present situation.
Conversely if a machine is to be replaced in two months time, it is pointless using a wear resistant
alloy that would increase working life by 12 months.
The function of a wear resistant alloy is to provide a more wear resistant surface to the type of
wear being experienced by the base metal. The wear resistant alloy can also often be applied in
such a manner as to physically entrap the wearing media thus giving a situation where the wearing
media is itself the wear protection overlay.
For the wear protective alloy to offer optimum protection, its selection must be based on the type
or wear and/or wearing media. The size of the wearing media will also affect the type of wear, for
example large rocks will produce more impact, and where applicable spacing of the pattern in
which the wear resistant alloy is applied.

11 SELECTING THE OPTIMUM WEAR RESISTANT SOLUTION
Assess mode of failure or wear type(s)
WHAT NEXT?
Establish the Facts
Properties of Deposit?
- thick
- work hardening
- heat resistant
What are the Service Conditions?
What is Causing the Wear?
What is the Desired Life?
What is the Present Solution?
Method of Application
- Stick
- Wire
- Powder
- Epoxy
- Compatibility?
Base Metal Considerations
- Susceptibility to cracking
- Compatible with wear facing solution
- Pre/post heat treatment?
- Dimensions/location of component.
What are the metallurgical implications of wear surfacing?
Welding Skills/Availability
- Process available
- Operator skill required
- Time for repair
Is the Solution Cost Effective?
Evaluate Prior Experience
- E+C CADB/TeroLink

12 METHODS OF DEPOSITION
For convenience, selection of wear facing type can be divided into two families.
(1) Fusion surfacing - e.g. arc welding.
(2) Non-fusion surfacing - e.g. ceramic sprayed coating.
Typical Characteristics Of Fusion Coatings
Process Dilution Thickness Deposition Rate Typical Uses
mm kg/hr
MMAW 15-30% 3 1-4 Quick, easy and local
repairs.
GTAW 5-10% 1.5-5 <2 High quality; low dilution
GMAW 15-30% 2-4 3-6 Faster than MMAW; no
slag.
FCAW 15-30% 2-4 3-8 Similar to GMAW but no
gas or flux required.
On-site use.
SAW 15-40% 3-5 5-30 Automated; heavy sections.
PTAW 2-10% 2-4 1-5 High quality; low dilution.
Automated.
Check compatibility of process with deposit requirements.
- base metal, etc.
- type of wear.

Typical Characteristics Of Non-Fusion Coatings
Process Dilution Thickness Deposition Rate Typical Uses
(mm) (kg/hr)
Flame Spray n/a <3 1.0-20 Zn/Al corrosion protection.
HVOF n/a <3 1.0-10 W2C/Co composites.
Arc Spray n/a
- Wire n/a 2-5 1.0-20 Shaft, journal reclamation.
- Plasma n/a 0.1-1.5 0.5-5 High quality, ceramics,
Aerospace industry.

13 WELDING PROCEDURAL GUIDELINES
Choose - Surfacing consumable
- Welding process
WHAT NEXT?
13.1 BASE METAL CONSIDERATIONS
Recognize potential problems such as:
HAZ embrittlement
Loss of strength hardness in Q&T steels
Reduction in corrosion resistance
Porosity generation from base metal chemistry
Contraction cracking
Consequence of fracture
Locate specification
Spark analysis
Component function
13.1.1 Weldability Factors
Other factors for consideration:
Cost
Weldability of base metal
Preheat
Postheat
Base metal properties

14 APPLICATION OF WEAR PROTECTIVE COATINGS
Once the wear resistant alloy type and process have been selected, see Section 7, the
method of application needs evaluation. Included in this is the preparation of the base
metal.
14.1 BASE METAL PREPARATION
The preparation will depend upon the condition of the base metal. If the component
has not previously been wear protected, only slight mechanical cleaning may be necessary.
Alternatively, previously cracked and damaged “hardfacing” may be present, which should
be removed by grinding, gouging, etc. depending upon extent. If the old hardfacing is left
on and covered, this may cause chipping of the deposit while in service.
All sharp corners and edges should be radiused, if that area is to be surfaced.
14.2 PREHEAT
This will depend upon the type, (e.g. never preheat 13% manganese steel), and
thickness of the base metal, i.e. if the part is large there is a greater heat sink effect
therefore more likelihood of hardening and crack susceptibility of the heat affected zone,
therefore preheat is advisable. Preheat will also depend upon available facilities and if a
highly alloyed buttering layer is used which reduces (although not usually eliminates) the
need for preheat.
Preheating carbon steels is usually based on the Carbon Equivalent
When to Preheat
CE Weldability Preheat Postheat
<0.45 Good Optional Optional
0.45-0.60 Fair 150-250o
C Preferable
>0.6 Care >250o
C Necessary

14.3 BUILD-UP
The extent of the build up will vary. In some instances where a lot of build up is
required, (e.g. earth moving tracks), grouser bars are welded onto the track as it is usually
easier, quicker, and more economical than rebuilding with arc welding (grouser bar may be
wear faced).
In other cases, build up layer(s) may be necessary prior to a layer(s) of the harder
abrasion resistant alloy, which should usually be limited to three layers maximum or in
some cases only one layer. If more than this is deposited cracking or spalling may occur.
A buttering layer of a highly alloyed electrode may sometimes be necessary with
harder more highly alloyed base metals to tolerate the dilution without cracking.
14.4 APPLICATION TECHNIQUE
See Section 10 - Wear Patterns.
14.5 COOLING PROCEDURE
If the base metal is hardenable the part should be cooled slowly after welding to
avoid cracking.
14.6 FINISHING
In certain wear systems, surface finish can affect wear life. With frictional/adhesive
wear the smoother the surface the better the wear life. With erosive wear, certain
investigators have reported an increase in wear life with smoother surface finishes. In
corrosive environments, a smooth surface eliminates the possibility of differential aeration
which results in accelerated attack.

15 WEAR PATTERNS AND PRODUCT SELECTION
It is often advantageous to apply a wear protective coating in a pattern or in a selected area
rather than covering the whole part with the alloy because:
(I) It can be deposited in less time.
(ii) A wear pattern can encourage entrapment of the wearing media and thus act
as a wear protective coating or it may encourage easy flow of material across the face of
the part.
and
(iii) It may encourage a self sharpening action, (see later).
Wear patterns are normally associated with abrasion, impact or abrasive/impact wear and
only these conditions will be considered.
15.1 WEAR PATTERNS FOR ABRASIVE AND IMPACT WEAR
The type of wear patterns used is dependent upon the type and size of the wearing media
and to some extent the base material.
The patterns must either prevent contact from the wearing media with the base metal by
itself providing the barrier, or it must encourage the wear media to become entrapped
inside the pattern thus preventing further contact with the base metal.
Let us consider the type of wearing media and then suggest possible wear patterns.
Large Particle Wear, i.e. rocks.
It is not possible to entrap rocks in a wear pattern to provide protection. Therefore we
must consider either complete coverage of the base metal to prevent contact, or lines of
wear facing material suitably spaced to prevent contact.

The recommended type of wear pattern is to deposit weld material resistant to impact (and
usually abrasion) in parallel lines to the flow of material. The rocks will then ride along
the weld deposit and will not come in contact with the base, i.e.
Direction
of
Flow
There are variations of this type of pattern to increase deposition coverage, e.g. “speed
dash”.
Direction
of
Flow
An important aspect is the correct spacing of the weld deposits. The lines of the weld
metal must be sufficiently close as to prevent the ingress of wear between them. The weld
deposit must also be resistant to impact and abrasion.

Fine Particle Wear, e.g. sand, earth, etc.
With this type of wear it is normally possible to employ a wear pattern to entrap the
wearing media. A recommended wear pattern is weld beads across the direction of flow of
material.
Direction
of
Flow
The spacing between the weld beads will be dependent upon the size of the wearing media
and the amount of moisture present in the material. A “wet” wearing media will compact
and become entrapped between the wear pattern more easily than a “dry” wearing media.
If the material does not become entrapped in between the wear pattern then the benefit of
the wear protective overlay is largely wasted. With this type of wear a material to resist
severe abrasion is required.
Large and Fine Particle Wear, e.g. rocks and sand.
This is the most commonly experienced wearing media in the open cast mining industry ⇒
a mixture of fine earth, sand, stone, etc. and varying sizes of rock.
The recommended wear pattern is a mixture of that needed to resist fine particle wear and
large particle wear, this being a “diamond” “waffle” or “crosshatch” configuration.
Direction
of
Flow
The fine wearing media should become entrapped in the pattern with only parts of the
wearfacing exposed.
THE SPACING MUST BE SMALL ENOUGH TO ENCOURAGE ENTRAPMENT

This technique is well highlighted in the picture below:
However, the wear pattern is often applied incorrectly as shown below:
The spacing of the pattern is too large and material has not become entrapped over the
whole surface, only in isolated areas. The solution here is to reapply a wear protective
coating but with a smaller spacing, e.g. 2 inch square “waffle”.

To protect the lip of the bucket it is recommended to completely cover with weld beads at
right angles to the direction of flow.
Direction
of
Flow
With a mixture of fine and large particle wearing media a weld protective overlay to resist
abrasion and impact is required.
“QUICK DOT” PATTERN
This is a method of covering a large surface fairly rapidly and with a minimum of heat
input (particularly advantageous with 12-14% manganese steel).
The “quick dot” pattern is usually used either to prevent contact of the wearing media or to
encourage entrapment of the wearing media by placing inside a “crosshatch” pattern. If
the quick dot pattern is used to prevent contact it must be placed closely together and not
as shown below:

BUCKET/DIGGER TEETH
Taking bucket/digger teeth as an example of a typical wear protective coating application,
it can be used to show the importance of not only which “wear pattern” to use but also
where the pattern should be applied.
With conditions of abrasion only, such as moving sand/earth it is generally accepted that a
wear pattern at right angles to the direction of flow is optimum and encourages entrapment
between the weld beads. Spacing must be close enough to encourage entrapment, typically
25-50cms.
SAND (Fine)
With impact conditions which occur when moving large rocks, a wear pattern parallel to
the direction of flow is recommended so that the rock rolls along the weld beads and does
not come in contact with the teeth. The spacing will depend upon the size of the rocks.

ROCKS (Large) COMBINATION
With conditions of fine and large particle wear a combination of the two patterns is
recommended, i.e. the diamond, cross-hatch or waffle pattern.
As shown on the above examples, the “quick-dot” pattern can be used in a variety of wear
conditions and is often easier and quicker to deposit particularly on vertical faces.
SHOULD ALL OF THE TOOTH BE COATED?
If the entire surface of the tooth is covered with wear protective coating, the nose of the
tooth will eventually wear on all faces resulting in a blunt tooth, reducing the working
efficiency.
ALL

If only the bottom surface of the tooth is coated the top will wear until there is insufficient
support for the base which will then break.
BOTTOM ONLY
The generally accepted technique is to coat only the Top and Sides, such that the base will
preferentially wear (but not excessively) and maintain the correct profile for its digging
action.
TOP SIDES
This self sharpening action can be applied to many industrial components working in
abrasion and abrasion/impact conditions.

15.2 WEAR PLATES AND GROUSER BARS
15.2.1 Wear Plates
This method of wear protection is sometimes suitable for covering large areas that are
exposed to abrasion and abrasion/impact.
There are two main types of wear plates:
(I) Weldable grades of wear plate - these can be welded into a suitable wear
pattern
(II) Welded wear resistant plates that are covered on one side with a wear
resistant overlay - has only to be welded in position (generally using low
hydrogen mild steel).
There are various factors to be taken into consideration when choosing between wear
plates and a weld protective overlay.
♦ Type of wear.
♦ Increased weight of the bucket/shovel, etc. using wear plates can cause more strain on
the equipment and decreased pay load.
♦ Adaptability, welded plates are usually large and not suitable for covering smaller areas
or radiused surfaces.
♦ Availability?
♦ Cost, welded wear plates are usually fairly expensive.
15.2.2 Grouser bars (e.g., bars for rebuilding worn
EARTH-MOVING TRACKS).
From the Manufacturers - The bars are normally made from medium carbon steel, e.g.
EN8 (similar to 1335, a 0.35%C, 1% Mn steel) and then induction hardened to create a
more abrasion resistant surface. New bars generally should not be welded as welding onto
the hardened layer can cause cracking.
Using Locally Available Steel - Use EN8 bar which is tough but weldable and then apply a
wear resistant alloy.

The EN8 bar should be welded to the track using an electrode designed for welding
dissimilar steels and then coated with a wear resistant alloy that is resistant to pressure and
abrasion, (see Sections 12 and 13).

16 SURFACING ALLOYS
16.1 CHROMIUM CARBIDE WEARFACING ALLOYS
Manual Chromium Carbide Arc Electrodes (MMAW)
Hardness Key Characteristics Features
57-60 Rc Fine to medium particle Can use contact welding
abrasion
57-60 Rc Abrasion/compression/impact Smooth deposit
Semi Automatic Welding (FCAW)
55-60 Rc High resistance to abrasion High hardness with
one pass
55-60 Rc Impact/abrasion/erosion Excellent weldability

16.2 WORK HARDENING ALLOYS (AUSTENITIC MANGANESE STEEL)
Manual Metal Arc Electrodes (MMAW)
10-15 Rc Austenitic deposit with Can be flame cut;
high crack resistance machinable
work hardens:
45-50 Rc
Semi Automatic Welding (FCAW)
20-25 Rc High resistance to cracking Machinable
work hardens:
45-50 Rc

16.3 IRON BASED BUILD UP AND WEARFACING ALLOYS
28-32 Rc Resistance to deformation Machinable and
and severe impact; build up economic and cushion layer
50-55 Rc Crack resistant Multi-pass build up
58-62 Rc For edge retention, high Heat treatable
temperature oxidation. Ultra
hard homogeneous deposits.
60-65 Rc Fine particle abrasion High hardness
with only one pass
58-62 Rc Abrasion/erosion Minimum dilution

16.4 TUNGSTEN CARBIDE WEARFACING ALLOYS
68-72 Rc Homogeneous deposit. High resistance to
High hardness in one pass. abrasion.
Torch Brazing/Welding Alloys
matrix: 200BHN Excellent cutting Low bonding temperature,
carbide: 89-91 Ra properties along with Good wettability
abrasion/impact
resistance
Extreme abrasion
self fluxed and
easily deposited

16.5 NI BASED WEARFACING ALLOYS
15-35Rc Impact, heat and corrosion Hot toughness
resistance
Torch Brazing/Welding Alloys
55-62 Rc Non magnetic deposits Low bonding
for wear due to friction and/or temperature; can
corrosion also be used with GTAW

17 GRADING OF WEAR RESISTANCE OF HARDFACING
ALLOYS
TYPE FEATURE TYPICAL APPLICATIONS
1. Tungsten Carbide
Composites
Utmost abrasion resistance
and moderate impact
resistance
Sand mixer blades; oil-well fishing tools; rock drills; tool
joints; dry cement pump screws; ripper tynes; pugmill
knives; ditcher teeth; coal cutter bits and picks; beaters;
post-hole auger teeth; churn drills, etc.
2. Chromium Carbide
Austenitic Irons
Excellent abrasion resistance.
High impact and oxidation
resistance.
Dredge buckets and lips; shovel and dragline teeth;
crusher jaws, mantles and rolls; bulldozer cutting edges
and end plates; muller pan tyres and pathways; impact
breaker blow bars; hammers; ball mill liners; tractor
grousers; agricultural implements; augers; rolling mill
guides; pump.
3. Martensitic Irons Excellent abrasion resistance.
Medium to fair impact and
oxidation resistance.
Similar to chromium carbide austenitic irons but not in
cases where high impact resistance is required.
4. Cobalt Base Alloys High to medium abrasion and
impact resistance. Excellent
hot hardness; corrosion and
oxidation resistance.
Hot and cold shear blades; metal and wood cutting tools;
dies; valves and valve seats; dishing; flanging; forming and
trimming dies; tap-hole augers; coke pusher shoes;
expeller worms; guillotine, sugarcane and mower blades;
pump sleeves and shafts; cams and tappets.
5. Nickel Base Alloys Medium abrasion and high
impact resistance.
Excellent hot hardness;
corrosion and oxidation
resistance.
Valve bodies, stems and seats; pump parts, flanges and
couplings; drawing, hot trimming and punching dies; acid-
resistant scrapers, etc.
6. Martensitic Steels Medium abrasion and impact
resistance.
Shovel track rollers, idlers and driving tumblers; tractor idler
wheels, rollers, track links and sprockets, etc.; rolling mill
rolls; wheel treads and tyres; clutch parts; guillotine blades;
punches; shears; dies, etc.
7. Pearlitic Steels Fair abrasion resistance and
high impact resistance.
Tractor idler wheels, rollers, track links and sprockets;
shovel track rollers, idlers and driving tumblers; bulldozer
arm trunnions; carbon steel rails; reclaiming of worn steel
parts (other than austenitic steels) prior to hardfacing.
8. Austenitic Steels
(Manganese and
Stainless)
Fair abrasion resistance.
Excellent impact resistance.
Work harden under some
conditions.
Dredge driving tumblers; austenitic manganese steel rail
points and crossings; digger teeth subject to extreme
shock; reclaiming of worn austenitic manganese steel
parts, e.g. crusher jaws, mantles, rolls, etc., prior to
hardfacing.

18 METHODS OF WEAR PROTECTION - SUMMARY
Essentially, wear surfacing involves placing a barrier between the wearing media and the
wearing part. There are various methods of doing this and by far the most common is the use
of surface modification.
When considering exposed working surfaces and the type of wear, the alternative solution is
to change the surface so making it more wear resistant. Basic methods of modifying the
surfaces include:
♦ By changing the metallurgical composition and/or characteristics, e.g. quenching and
tempering of hardenable alloys.
♦ By changing the condition of the wearing media.
♦ Physical separation by the use of “wear plates”.
♦ Physical separation by the use of “wear protective overlays”.
IMPORTANT
For optimum increase in wear life and economic benefit the wear protection should be
undertaken in most cases when the part is new or after slight wear.

19 REFERENCES
Below is a brief list of additional sources of information that you may find valuable .
Metallurgy for Engineers
Physical Metallurgy for Engineers
Weldability of Steels
Welding Handbook, 7th
Ed., Vol. 4, AWS
Metals Handbook, 10th
Ed., Vol. 6, ASM
WELDING AND COATING METALLURGY2 69of 69
22 December, 2002;20:30

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